High strength alloys and methods for making same

ABSTRACT

A family of extremely fine-grained alloys are used to make coatings or free-standing bodies having desirable properties for use as a heat-resistant and wear-resistant material. In an illustrative embodiment, the alloys are comprised of a multiplicity of alternate, microcrystalline or nanocrystalline films of tungsten metal and tungsten compound. The tungsten compound film may be comprised of a tungsten carbide or a tungsten boride. The tungsten films are the primary films. Their desirable characteristics, in addition to their very fine crystalline habit, per se, are the high strength, high hardness, high resilience, and high fracture energy which these fine crystallites foster. They may be manufactured by a chemical vapor deposition process in which reactive gas flows are rapidly switched to produce alternate films with abrupt hetero-junctions and thereby to produce the useful micro-crystalline habit. The unique synthesis method allows effective control of critical flaw size. The structure is such that the primary films may be made sufficiently thick so as to assure some desirable ductile behavior, but sufficiently thin so as to have high yield strength by dint of their microcrystalline size, and as to limit the size of any flaws. The secondary films are made of enough thickness to prevent the epitaxial growth from one primary film to the next-deposited primary film and thin enough so that they can not contain a flaw of critical size. In addition, the exterior surface of any body made by this method may have a sufficiently smooth surface that the strength of the body is determined by the bulk properties of the material and not by surface flaws.

RELATED APPLICATIONS

The present application claims priority from U.S. provisionalapplication Ser. No. 60/246,637 filed Nov. 2, 2000 and is incorporatedby reference herein.

FIELD OF THE INVENTION

This invention relates to heat-resistant, and wear-resistant alloysuseful for coatings or free-standing bodies having favorablecombinations of strength, hardness and/or toughness. It also relates tomethods for making the alloys. More specifically, the invention relatesto chemical vapor deposition processes and products therefrom, whichhave unique, substantially improved physical and mechanicalcharacteristics.

BACKGROUND OF THE INVENTION

Group VIB transition metals include, for purposes of this application,tungsten, molybdenum and chromium. The Group VIB, transition metalelements, such as tungsten, molybdenum, and chromium, havecharacteristics that allow their incorporation into some new,high-performance alloys. Their high stiffness suggests that they haveintrinsic high strength. This indicates that they should have highfracture energy and high specific resilience. It also suggests that theyare capable of being made into hard and wear resistant alloys. They havehigh melting temperatures, as well. Unfortunately, the potentiallysuperior mechanical properties of these materials are seldom realizedbecause of their lack of toughness.

They are all used as pure metals and as important alloying species withbase metals. As pure metals, or as the major species in alloys, tungstenand molybdenum are more important than chromium as structural materials.Chromium is used more frequently as a coating.

These Group VIB transition metals, such as tungsten, are usedindustrially as pure metals, sometimes containing small quantities of afinely divided dispersant; as an alloy with other high melting metals;or as a pure metal cemented into a body with small quantities of a lowermelting metal matrix; or as a carbide, either pure or alloyed, cementedwith a similar lower melting metal matrix. They are also used as adilute alloying species in high strength and high hardness base metalalloys.

Among the most important applications of tungsten, for example, are forresistance wire as in lamp bulbs and vacuum tubes, extremely smallconductors in microprocessors, x-ray targets, so-called heavy metalalloys, and cemented carbide tool and wear parts. The wire and x-raytarget uses take advantage of tungsten's high melting temperature; themicroprocessor use of its electrical conductivity and thermal expansioncoefficient; the heavy metal alloys of its high specific gravity; andthe cemented carbides of the hardness and wear resistance of itsmonocarbide.

In most instances, it is important for these Group VIB transition metalsto have the highest strength and toughness, consistent with themaintenance of its other important properties.

Fine tungsten wire, for example, after the large amount of mechanicalwork which goes into its manufacture, exhibits high strength. Bulk metalparts of tungsten are usually much weaker, however. In all but a fewinstances, e.g., the fine wire, tungsten parts suffer from lack oftoughness. Even the wire soon loses both strength and ductility onheating due to the work being a high driving force forre-crystallization and grain growth. The brittleness of x-ray targetsand other larger bodies has been avoided, at considerable increase incost, by the addition of the rare metal, rhenium, as an alloying speciesin quantities as high as twenty-five percent.

The heavy metal and cemented carbide parts rely on another approach toachieve acceptable toughness. They are made by pressing and sintering amixture of pure metal powder, or of carbide powder, with alower-melting, more ductile, base metal. The tungsten or tungstencarbide is thereby cemented by the small quantities of the ductile basemetal.

Properties of the final product are achieved by the judicious selectionof the matrix metal composition, the size of the metal powders, or thesize and composition of the carbide powders. For many applications oftungsten and for most applications of tungsten carbide thebase-metal-cemented, these pseudo-alloys are the only practicalsolutions. There are many instances, however, where the incorporation ofthe softer, lower-melting, less-stiff, and less corrosion-resistantcement substantially degrades the usefulness of the bodies. Puretungsten, or alloys of tungsten with strengthening or hardening specieswhich would not use such cement would be much more useful.

With regard to metals and other materials in general, it has been wellknown to materials engineers and scientists that refinement of thecrystal habit of bodies increases yield strength, and hardness. Sinceancient days mechanical working to reduce their grain size hasstrengthened metal parts. With more sophisticated understanding, theso-called Hall-Petch relationship has become generally accepted. Thisrelationship teaches that the yield strength of materials variesinversely with the reciprocal of the square root of the grain size. In amore recent publication, Jundal and Armstrong (see Trans. AIME 1969 vol.245, pg. 625) reported that the Hall-Petch relationship could beextended to treat the increase in material hardness with grain sizereduction as well as yield strength. Additional verification, for thecase of the hardness of tungsten, comes from Vashi, et al. (seeMetallurgical Trans., Vol. 1, June 1970, pg. 1769-1771). (The entirecontents of all publications and patents mentioned anywhere in thisdisclosure are hereby incorporated by reference.)

Within the last decade, research has demonstrated that the dramaticeffects on properties can be extended in materials of much finer grainrefinement than had been earlier possible. Progress in the manufactureof cemented tungsten carbide cutting tool materials discussed above is aparticularly good example of such improvement. Two decades ago the mostmodern of these cemented carbides had WC crystallite sizes no smallerthan about two microns. Today, they are made quite regularly,commercially, with 0.4 micron (400 nm) crystals; and even smaller, on anexperimental basis. This has resulted in superior products from thepoint of view of strength and wear resistance.

This reduction in grain size is not accomplished without difficulty.There are practical limits to the fineness of powders which may be usedin the pressing and sintering process. Very small powders have long beenconsidered explosion and worker-ingestion hazards. Even moreimportantly, these powders tend to agglomerate in handling, therebypreventing the formation of a final product with a crystal refinement assmall as might be desired.

Advances to reduce the agglomeration problems have been claimed to beeffected by the use of a spray-reaction process from salts of tungstenand the matrix metal with subsequent gas-phase carburization. Thisprocess is described in U.S. Pat. Nos. 5,230,729 and 5,352,269. Further,however, even after these very fine powders have been pressedsuccessfully to a so-called green body, there is a tendency toward graingrowth upon sintering, although efforts have been made to alloy thecementing metals to allow lower temperature processing and to minimizethis grain growth. This approach is described in U.S. Pat. No.5,841,044.

For reasons which have not been totally explained, none ofsub-micron-size or nanostructure cemented carbides, except those withgrain sizes above about 0.4 μm, or even above 0.8 μm, has shownsufficiently good toughness to be generally accepted commercially.

In the materials science arena, however, investigators have becomeincreasingly anxious to investigate the effects of nano-technology.Nano-technology is usually defined as dealing in microcrystalline sizesbelow 0.1 μm (100 nm). Because of the aforementioned limitations, andbecause they need only small samples, they have chosen to use depositiontechniques to make their research samples. Deposition is an attractiveway to make extremely fine-grain materials since the crystallites of thematerials of interest may be grown and consolidated, simultaneously, attemperatures which are low relative to their fusion temperatures, oreven to their sintering temperatures. These bodies made by variousdeposition methods, therefore, need not be limited as to their coarsecrystalline habit, as in casting; or as to agglomeration, or graingrowth, as in powder pressing and sintering. Properly manipulated, theycan be consolidated to virtually full density, quite free of internalvoids and defects.

Both electrochemical deposition (ECD) or physical vapor deposition (PVD)techniques have been used by these scientists to make such samples fortheir scientific investigations. In the present application, physicalvapor deposition refers to any of the group of similar methods,including evaporation, reactive evaporation, sputtering, reactivesputtering, and ion-plating. Such efforts are described in papers byMenezes and Anderson; J. Electrochemical Soc. 137, 440 (1990) and Chuand Bamett; J. Appl. Phys., Vol 77, No. 9, 1 May 1995. The samples havebeen useful to investigate the achievable improvement in properties frommaterials with grain refinement smaller than 0.25 micron (<250nanometers). Small-scale samples have been made and tested. They haveusually been made of a multiplicity of thin layers. Films withcrystallite sizes well below 100 nm (even below 10 nm) have beensuccessfully synthesized. These techniques, however, approach theobjects of high performance materials in a very different way from thoseof the invention. They did not involve the strengthening and hardeningof a metal with some intrinsic toughness, but rather an investigation ofwhat happens when the grains of an intrinsically brittle material arerefined.

It has been determined that much greater hardness can be achieved insuch materials by the aforementioned techniques. However, improvementsin strength or toughness have not been generally measured and reported.

The chemical vapor deposition (CVD) process would be more appropriatefor the manufacture of industrial parts of the materials of interestthan ECD or PVD. CVD, although requiring processing temperatures higherthan either ECD or PVD, can still be processed well below the requiredfusion temperatures or sintering temperatures for the materials ofinterest. In the present application chemical vapor deposition is meantto include both simple thermally-activated CVD as well asplasma-assisted CVD. Since the control of CVD is more difficult thanthat of either ECD or PVD, it has been used very sparingly for any kindof nanotechnology research and hardly at all for any commercialmanufacture of such fine-grain materials.

There are three notable exceptions. The most significant one isdescribed in U.S. Pat. No. 4,162,345 to Holzl ('345). Two decades ago,the inventor, Holzl (one of the co-inventors of the current invention)taught, in the '345 patent, that materials made by a then-uniquevariation of the CVD process could be made to demonstrate a usefulcombination of strength and hardness such as to provide excellent wearresistance. The material could be described as an early version of whatis currently being called a nanostructure.

The second is a research program conducted at Stevens Institute ofTechnology by Eroglu and Gallois in which thin nanostructure TiN/TiCcoatings were investigated (see “Design and Chemical Vapor Deposition ofGraded TiN/TiC Coatings”; Surface and Coatings Technology 49, 275(1991)). Like the Chu and Barnett work, cited above, these investigatorstook a different approach than that of the invention. They were alsoinvestigating the refinement of normally brittle materials.

The third is a wear-resistant coating for cemented carbide tools whichhas been offered commercially since late 1998 by Widia Valenite. Thatcompany introduced a thin, nanocrystalline coating for cutting toolsmade by what is called multilayer CVD (MLCVD). They report improved wearlife for certain cutting applications and claim that crack formationthrough the entire thickness of the coating is minimized by themultilayer configuration. The reported coating is comprised ofconventional, brittle, coating materials, titanium nitride and titaniumcarbonitride. No improvement in strength was reported, or should havebeen expected from this work.

The background art closest to the current invention is the referencedwork of the '345 patent. Most of the microcrystallites in the Holzlmaterial were in the order of 50-100 nm, but it contained some that wereas much as ten times larger. The material was actually used for certainimportant valve trim in the NASA space program. Unfortunately, themethod of the '345 patent could not be reproduced with acceptablereliability and was extremely expensive. The irregularity of thecrystallite size was a major problem which was never adequately solved.The process was subsequently discarded as unacceptable for industrialuse.

However, there is ample reason to continue to be interested in CVD as aprocess for making nanostructural parts. Electrochemical depositionwhile totally acceptable for the common metals is practically uselessfor making refractory materials such as tungsten, its alloys, orcompounds. PVD can be used to make common metals and compounds at veryhigh rates, but for refractory metals and alloys, deposition rates areunacceptably low. Chemical vapor deposition, on the other hand, cansynthesize such refractory metals and ceramics at very acceptablecommercial production rates. CVD is superior to ECD and PVD asmanufacturing processes in other ways, as well. Principal among them areits excellent throwing power and its ability to make materials of higherand essentially full density, virtually free of internal voids. The useof CVD to produce high melting and chemical and wear-resistant metalsand ceramics is well-known.

Processes for making free-standing shapes of the so-called refractorymetals and alloys have been known for decades. For example, puretungsten tubing has been made commercially by depositing the metal on amandrel from which it is then removed. Parts of good purity exhibit aVicker's hardness of about 4 GPa. Utilizing the method of Cahoon et al.(see J. B. Cahoon, W. H. Broughton, and A. R. Kutzak, MetallurgicalTransactions, vol 2, pp. 1979-1983, 1971), which teaches that the yieldstress of a material that is fully strain-hardened is approximatelyequal to one third the Vicker's hardness, a maximum yield stress forhigh purity, CVD tungsten of 1300 MPa would be predicted. This value isan upper anticipated limit as the CVD tungsten is not fullystrain-hardened. In practice, maximum values of 900 MPa can be obtainedand the material displays limited ductility at room temperature. Thecolumnar growth of the CVD tungsten produces near-continuous grainboundaries, which act as a volumetric flaw within the material. Thisstructure leads to brittle failure of the tungsten at room temperaturewith strengths closer to 300 MPa for larger-grained, free-standingshapes. The corresponding low fracture toughness limits their utility.Reducing grain size would be expected to increase the strength. Onetechnique for so doing is the lowering of the deposition temperature.The process then suffers from reduced deposition rates. A secondtechnique involves mechanical burnishing the work piece duringdeposition (see L. W. Roberts; Proc., Sixth Plansee Seminar, Reutte,Austria, 1967; pp 881-884). This is mechanistically difficult on all butthe simplest of work pieces. The highest strengths achievable by eitherof these techniques is, at most, about 900 MPa. All of the strengthvalues cited above, and those that follow in this application, areflexural strengths, measured in 3-point bending with 24 mm diameterround rods.

Brittle materials like the refractory carbides, nitrides, borides andsuicides are also conveniently made by CVD. CVD-synthesized tungstencarbide, having a hardness above 20 GPa, is not likely to have astrength of greater than 70 MPa and is essentially useless as anythingbut adherent thin coatings.

CVD has been used for years to produce thin films, such asoxidation-resistant coatings for high temperature metals andwear-resultant coatings for a wide variety of cutting tools. In the caseof such thin coatings as these, c.a., 4-8 μm, the deposits are notrequired to have significant strength since their structural integrityis derived from the substrate upon which they are adherently deposited.Thin coatings of this kind which also were extremely fine-grained aredescribed in U.S. Pat. No. 4,427,445 to Holzl.

The process described in the '345 patent and other related patents byHolzl is the most notable claim of using CVD as a means of producingmetals, semi-metals, or refractory compounds having an uniquecombination of high-strength and excellent fracture toughness,especially in materials of high hardness.

Following the teachings of the aforementioned Hall-Petch relationship,there was reason to believe that the characteristic of the materialsdescribed in the '345 patent which caused them to have such uniqueproperties was their extreme grain refinement, c.a., 50-100 nm. Theirhigh hardness was attributed, at least in part, to their content oftungsten carbides. There is also ample reason to believe that thevariability which was experienced in the products made by the method ofthe '345 patent was due to the presence of some irregular, larger grainsin the structure.

In the specification of the '345 patent, Holzl postulated that theformation of the microcrystalline grain structure was a result of areaction off of the surface of the substrate to form a liquidintermediate product which was subsequently reacted to form a secondliquid intermediate product, which is deposited on the substrate,thence, rapidly, to be reacted to form the desired solid phase. In thisrespect, this process might be the equivalent of very rapid quenching ofa metal from the melt which has been used to cast extremely fine-grainmaterials. Such a sequence of events probably occurred but, was not, inand of itself, sufficient to fully explain the results.

Holzl also postulated that the observed layered structure was caused byoscillating turbulence in flow of the fog or halo off of the substrate.This is now believed, based on the investigations of this invention, tohave been an absolutely essential factor in the described depositionbehavior.

The near impossibility of causing this oscillation to occur in a totallypredictable way was most likely the fundamental cause for the processbeing non-reproducible and discarded as not commercially practical. Eachtime that the size or shape of the deposition reactor was changed andeach time that the size or configuration of the work piece(s) waschanged, an entirely new set of deposition conditions needed to bedetermined to establish this oscillating turbulence properly.

The process was so sensitive that even minor changes in the positioningof the work pieces in the reactor could cause failure of the processingruns. The layered structure was simply not acceptably uniform in itsfrequency and thickness of its layers.

The material made according to the method of the '345 patent wasconsidered to be an alternate and improved method to the powdermetallurgy of cemented carbides for the making of hard metal parts fortool and wear applications. It was considered to be superior to cementedcarbides because it eliminated some of their deficiencies. In manycases, the wear resistance of cemented carbides is dictated more by theperformance of the cement than by the hard particles and is therebylimited. In short, wear occurs frequently by the failure of the cementallowing the hard particles to be removed from the body without theparticles, themselves, actually fracturing or wearing.

This behavior of cemented carbides can be compared with that of otherwear materials like tool steels. Tool steels, although they contain twoor more phases, wear like a homogenous material, not like a mixture.They also have greater toughness than any other materials of equivalenthardness.

If tool steels could be made as hard as the cemented carbides, theywould be much preferred. The same statement could be made about cast,hard nickel or cobalt alloys versus the cemented carbides. The maximumhardness of tool steels or the cast hard alloys, however, is typicallyonly about one half the maximum hardness of the cemented carbides; towit, ˜7-9 GPa Vickers Hardness Number (HV) as compared to ˜11-22 GPa.For this reason they are disqualified from many applications for whichcemented carbides can be used. An additional advantage of the cementedcarbides over tool steels is, of course, their ability to maintain theirstrength and hardness at the high temperatures generated within the toolmaterial in certain machining operations.

Note that as included in this disclosure, HV is used to denote VickersHardness Number as measured with a 500 g or 1000 g weight on a ShimadzuMicrohardness Tester, unless otherwise cited.

SUMMARY OF THE INVENTION

The current invention makes metal alloys which, in many respects, aresimilar to tool steels. It does not attempt to make brittle,ceramic-like materials stronger and tougher, as is done in cementedcarbides, but rather to make a more ductile, metallic material strongerand harder, as is done in tool steels. The invention utilizes Group VIBtransition metals, such as tungsten, molybdenum, or chromium.

In an illustrative embodiment of the current invention, tungsten isutilized as the primary material. However, it will be appreciated bythose of ordinary skill in the art that other Group VIB transitionmetals, such as molybdenum or chromium, could also be used and stillremain within the scope of the current invention. In the illustrativeembodiment of the current invention, the tungsten alloy is significantlytougher than many other forms of tungsten due to the near brittle natureof tungsten at room temperature.

In comparison to tool steels, the invention makes use of materials suchas tungsten as the major constituent, instead of iron. Intrinsically,tungsten is stronger than iron; this is because it is generally acceptedthat, for a given crystallite size, the attainable maximum yieldstrength varies directly as the square root of the ratios of thestiffness of any two materials. Tungsten having almost exactly twice thestiffness of iron may be expected to have an intrinsic strength of about1.4 times that of iron. The greater stiffness of the tungsten, per se,can be important in many applications.

In addition, tungsten, and tungsten alloys, with their high meltingtemperatures, maintain their strength and hardness at highertemperatures than do iron or iron alloys. Although the materials of thepresent invention are a mixture of two species (like tool steels orcemented carbides), the crystallite size of each of the species in theinvention alloys is so small that they, like the tool steels, and unlikethe carbides, act like a homogeneous alloys. In the followingdescription, they are, therefore, referred to as alloys.

Accordingly, it is an objective of the present invention to provide animproved method for producing alloys of Group VIB transition metalswhich have properties superior to alloys of such metals which arecurrently commercially available. These metal alloys are morecontrollable and more able to be tailored to desired properties ascompared to those described in the '345 patent, and the mechanism ofstrengthening and toughening is different from that purported in the'345 patent. The improved method allows for totally acceptablereproducibility, is capable of being scaled up to larger quantities ofproduct, can accommodate products of different sizes and configuration,and is inexpensive to operate.

It is an objective of the invention to be able reproducibly to makecoatings and free-standing parts of such metal alloys, which would haveunusual and desirable combination(s) of strength and fracture toughnessfor any applications in which conventional, lower performance alloysmight be used and exhibit significantly greater resistance torecrystallization than alloys whose fine-grain size is derived fromextensive mechanical working.

It is an objective of the invention to provide an alloy of adherentlayers wherein intermediate layers prevent epitaxial growth betweenadjacent layers.

It is an objective of the invention to provide alloys in which both thedeficiency of low strength and that of low toughness are avoided.

It is an additional objective of the invention to make alloys ascoatings or as free-standing bodies, which have such unusual combinationof strength, fracture toughness, and hardness that they provide longeruseful life as tools or other wear parts.

It is an additional objective of the invention to make alloys ascoatings or as free-standing bodies, which are able to be made directlyto include, or finished to include, surfaces of great smoothness byvirtue of which the bulk properties of the alloy can be achieved withoutsuffering from failure due to surface defects, and which, in addition,have utility for reflecting or low-sliding-friction surfaces.

It is an additional objective of the invention to make alloys which arecapable of being finished with very keen edges or very smooth surfacesrequired for many tool and wear parts.

Other objects and advantages of the present invention will become moreapparent to those skilled in the art from the following description,taken in conjunction with the accompanying drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic drawing of a type of CVD equipment used for theinvention.

FIG. 2 is a low magnification photomicrograph of a conventional tungstenchemical vapor deposit, showing coarse columnar grains. It is theintergranular cracking between these large grains that causes theweakness and poor toughness of such metals. It was necessary to use alow magnification to illustrate the structure because of the coarsegrains.

FIG. 3 is a medium magnification photomicrograph showing the uniformityof the primary and secondary film structure made possible by thisinvention. Higher magnification had to be used in this figure because ofthe very fine grains.

FIG. 4 is a higher magnification photomicrograph of a fine-grained,tungsten alloy of the invention. Note that the grain sizes are in therange of 150-200 nm, about in the middle of the range of typical grainsizes for the alloys of the invention.

DETAILED DESCRIPTION OF PREFERRED EMBODIMENTS OF THE PRESENT INVENTION

The nature of the alloys which are the subject of this invention is bestrevealed by the manner by which they are manufactured. An illustrativemethod utilizes CVD equipment of a generally accepted design, like thatshown in FIG. 1.

In a first illustrative embodiment of the method, a volatile halide oftungsten is supplied as a gas from a separate container, or is generatedin situ by reaction of the elements. It is clear that the volatilehalide of tungsten may be chosen from the group comprised of fluorides,chlorides, bromides or iodides.

The highest valence state fluoride, WF₆, would be an attractive choice.It is readily available commercially, is conveniently packaged, and hassufficiently high vapor pressure to allow its easy dispensing into a CVDsystem. In fact, it is the most commonly used precursor for tungstendeposition by CVD. Unfortunately, because of modern considerations ofenvironmental hazards, fluorine and any volatile fluorides are now veryexpensive; for example, WF₆ is so expensive as to virtually disqualifyit for most industrial uses except for premium applications such asmicroprocessors and the like. Tungsten chlorides, on the other hand, areinexpensive, have acceptable chemical thermodynamic characteristics andpresent minimal environmental problems in the deposition process, perse, or in any recycle or waste disposal processes indicated. There arethree sufficiently volatile tungsten chlorides: WCl₆, WCl₅, and WCl₄.

The tungsten bromides and iodides are clearly acceptable precursors forthe deposition process, but appear to offer no advantages over thechlorides and are more expensive. While this is also true formolybdenum, a volatile iodide is the most effective precursor for thedeposition of chromium.

The tungsten halide may be dispensed from an external container, usingheated lines and an inert carrier gas, if required, to assist in thetransport. Such heating and carrier gas are not required for veryvolatile tungsten hexafluoride. They are for the more commerciallydesirable tungsten chloride precursor, however.

FIG. 1 shows the more convenient technique of making the tungstenchloride in situ. Chlorine gas is metered through a heated bed oftungsten chips where it combines to form the necessary tungsten chlorideto transport the tungsten. Different kinds of chips may be used. Theterm chips is used to distinguish the feed material from powder or largechunks of metal, both of which are not satisfactory. The powder has thenecessary large surface area but can pack too tightly, thereby notallowing the gas to pass uniformly through the bed. It is also likely tobe entrained by the flowing gas. Large chunks have too little surfacearea.

Tungsten scrap can be very useful. Machining turnings or chopped wireare inexpensive and in good supply. Virgin powder can be pelletized andmade to work successfully.

Bed temperatures of between 600-900° C. are required; depending on thetotal surface area of the chips used and the flows of chlorine. Lowertemperatures are suitable for the low flows which would be used for thelow deposition rates such as in making thin coatings, but do not providefor sufficiently rapid reaction with commercially-required flows formassive coatings or free-standing bodies. Excessively high temperaturestend to drive the product to a non-volatile lower chloride.

Using this technique eliminates the necessity for using heated feedlines or a carrier gas or both, and minimizes the corrosion andcontamination problems of such a heated system.

To make a chemical vapor deposit of tungsten metal, the chloride flowsfrom the bottom of the chip bed into the reaction zone of a gas-tightvessel where it is made to encounter heated workpiece substrates. Thetemperature to which the workpiece substrates are heated depends onwhich tungsten halide is used and the quantity of hydrogen used as areductant. In the absence of hydrogen, a deposition temperature above120° C., would be required. With a hydrogen reductant and a tungstenchloride precursor, the temperature is selected to be between about 400°C. and 1100° C.

The workpiece can be heated by any of the standard techniques used forconventional CVD, such as the use of an internal heater, radiation froma hot wall or by induction. Lower temperatures are used for a lowdeposition rate, as, for example, in the formation of thin coatings. Incontrast, high temperatures are used for high deposition rates as arerequired for commercial acceptability in the making of massive coatingsor free-standing parts.

The volatile metal halide is decomposed by heat and the action of thereducing gas to form the desired deposit on the workpiece. Hydrogen, asshown in FIG. 1, may be utilized as a reducing agent.

The continuous flow of reactant products is continuously removed fromthe reaction zone by a vacuum pump, designed to allow operation atsub-atmospheric pressures to control the deposition rate and thecrystallite size, and to minimize any condensation or desublimation ofany of the reactant or product gases. As noted above, the temperature ofthe workpiece also is used to control the deposition rate.

Conventional CVD of tungsten and of compounds of tungsten are quitestraightforward. In the method of the invention, the deposit is not madeby a conventional CVD process, however. It does, however, produce thenear theoretical density and freedom from large flaws that arecharacteristic of CVD.

The process is started by the deposition of a very thin tungsten film onthe substrate. The tungsten film of the invention is typically about tento 1000 nm thick. The very thinness of the film additionally assuresfreedom from large flaws in the deposit. Such internal flaws are avoidedby eliminating both porosity and large grain boundaries. For purposes ofthe following discussion the tungsten film is called a primary film.

After the extremely short duration of deposition required to produce athin primary film (usually only a few seconds), an additional gas isconducted into the reaction zone so as to cause formation of a thin filmof a compound of tungsten, e.g., a carbide, boride, or silicide. Thisadditional gas is shown in FIG. 1 being injected with the hydrogen, butit may be injected at any point upstream of the workpiece(s). Theadditional gas is injected in short pulses, controlled by the timervalve shown in FIG. 1. For purposes of this description, this film willbe called a secondary film.

The CVD process, as regularly practiced, produces first deposits made onan indifferent substrate that are extremely fine-grained, unoriented,and, slower growing. After a period of time, however, the growth becomesoriented and faster.

In a method of the invention, the growth of each primarymicrocrystalline film is stopped, therefore, by the secondarymicrocrystalline film, before the coarser, oriented, crystallites begintheir growth. The material of the secondary film must be of sufficientlydifferent crystal habit from the body-centered cubic habit of theprimary film metal and of sufficient thickness as to interrupt thisoriented, rapid growth of the crystallites of the primary film. Thesecondary film must also be inactive with the primary film. In otherwords, the primary and secondary films must not readily react chemicallywith each other nor significantly dissolve in each other. The avoidanceof this rapid growth of the crystallites of the primary film isnecessary to prevent coarsening of these crystallites from theirinitial, very-fine habit and to control the thickness of any primaryfilm as necessary to effect the method of the invention.

FIG. 2 shows how the crystallites continue to grow in a columnar habitif they are not interrupted. This is shown at a low magnificationbecause the crystallites are so large.

FIGS. 3 and 4 show how the secondary film effectively interrupts thiscolumnar growth and maintains grain refinement on a nanocrystallinescale. These are shown at a much higher magnification.

A secondary film must also be very thin. This is because, first, itsrapid growth and coarsening must also be prevented, and second, controlof its thickness is a crucial part of the invention. Its thinness avoidsthe formation of flaws in the secondary film, itself. It can be seenthat one of the unique aspects of the invention, as compared withconventional CVD, is the high-speed, controlled switching of theprecursor gases.

In the illustrative embodiment, the secondary film is comprised oftungsten combined chemically with a another element to form a compoundof tungsten. In this embodiment, carbon and boron are the preferredelements. They have been determined to be acceptable in combining withtungsten to form secondary films which effectively perform the functionof arresting the growth of the primary film crystallites. The word formis used in the preceding sentence since it is intended to refer eitherto making the compound of tungsten by some conversion of thepreviously-deposited primary tungsten film or totally by deposit fromthe gaseous species. The words form and deposit are used interchangeablythroughout this application to describe the making of the secondaryfilms.

In the making of the material of the '345 patent, it was thought to beimportant that the second phase be a hard material. The high hardness ofsome of the embodiments of the alloys of the invention is believed to bedominated by the grain refinement, however. Although high hardness ofthe secondary films can clearly contribute to the hardness of theoverall material, secondary films of lower hardness can clearly makeuseful alloys. It would be obvious, however, to someone of skill in theart that whereas the material of the secondary films need not be hard toeffect the necessary function of stopping the growth of the primarygrains and preventing epitaxy, there would be advantage in its beinghard. These compounds comprising the secondary film in this embodimentcan properly be called hard metal compounds. In this description of theinvention we are following the definition of hard metals presented inthe well-known text, Refractory Hard Metals; P. Swartzkopf and R.Keiffer; The MacMillan Co., New York, 1953; Chapter 1.

The crystal structure of the secondary film composition does not appearto be of profound importance. For example, there are three reportedcarbides of tungsten. The uncommon W₃C is formed only at relatively lowdeposition temperatures; below about 600° C. It can be expected to be auseful secondary film for thin coatings deposited below thattemperature. At higher deposition temperatures; as in the deposition ofmassive coatings or free-standing bodies, both WC and W₂C have beenobserved. For example, when propane is used as the precursor for thesecondary film, the composition has been determined, by x-raydiffraction to be W2C. On the other hand, when propylene, which has ahigher carbon activity, was used as a precursor, the composition of thesecondary layer was determined to be WC. Either seems effective ininterrupting epitaxial growth of primary film crystallites. The twocarbides are of similar hardness and appear to have a similar slope inthe Hall-Petch relationship.

In the case of the tungsten borides, three compounds, W₂B, WB, and W₂B₅,are known to exist. Available thermochemical data would suggest that thelast would be unlikely to be made with any convenient depositionconditions. The behavior of W₂B and WB are expected to be the same asthe carbides as far as behavior as secondary films is concerned. Atungsten boride secondary film has been demonstrated, experimentally, tobe effective in interrupting the epitaxial growth of tungsten. However,the crystal structure could not be identified by x-ray diffraction dueto the small fraction of boride present in the alloy.

The disilicide of tungsten would appear, from thermochemical data to bethe only silicide of consideration as a secondary film.

There are specific reasons for having the secondary film contain atungsten compound. The reasons for this selection are discussed below.

When the secondary film is being formed, the flow of thetungsten-halide-containing primary film precursor gases is not turnedoff. The reactants for the secondary films are simply added to theformer. Suitable additive reactant gases for a carbide, boride, orsilicide secondary film are, for example, a volatile hydrocarbon or agas like carbon monoxide for a tungsten carbide secondary layer, or avolatile boron halide for a tungsten boride, or a volatile halide orhydride of silicon.

Many more compositions for the secondary film are possible: othercompounds of the primary metal, even including intermetallics, othermetals, semi-metals, other metal compounds, or semi-metal compounds.There are two considerations when using these other compositions as thesecondary film. The first would be to assure that abrupt heterojunctionsare effected between the primary and secondary films. This means thatany species used to make the secondary film preferably has limitedreactivity with, or limited solubility in, the primary film metal at theconsolidation and use temperatures of the alloy, as is the case with theprimary metal carbides, borides, and silicides.

There is also the matter of the possible necessity that turning off theflow of the primary-film deposition precursor might be required. Itcertainly would be possible to make the secondary films from othercompounds of the primary metal, even including intermetallics. However,to make the film of a different metal or of a semi-metal compound could,in most cases, require switching of the primary precursor. An exceptionto this switching requirement for a semi-metal compound secondary filmmight be silicon carbide since it can be formed by decomposing a singlegas in the presence of the precursor gases for the primary films.

This switching of the primary film precursor is considered to be lessconvenient from a commercial point of view. In order to insure that noneof the primary precursor remained in the deposition chamber, anadditional inert gas purge cycle might need to be included between eachthin film deposition, which would greatly increase, perhaps even double,the overall processing time for a deposition run. In any event, aprocess for depositing an alternative secondary film which involvesswitching between precursor gases is within the scope of the presentinvention.

Such a process for depositing an alternate type secondary film couldinclude the following steps: (1) Turning on the primary depositionprecursor and allowing it to flow for a selected deposition time; (2)Turning off the primary precursor; (3) Turning on an inert gas for atime sufficient to purge all of the former gas out of the chamber; (4)Turning on a secondary film deposition precursor(s) and allowing it toflow for a selected deposition time; (5) Turning on an inert gas for atime sufficient to purge all of the former gas out of the chamber; and(6) Repeating the cycle until a sufficient deposition thickness isattained.

Returning to the first illustrative embodiment of the process or thepresent invention, the important selection criteria for the compositionof the secondary films are:

-   -   1) that only a single additive gas be needed;    -   2) that this: single gas can be made to react with the tungsten        precursor gas(es) to form the solid secondary film at conditions        essentially identical to those used to deposit the tungsten in        the primary film, and is not so unstable as to react        homogeneously in the gas stream to form soot (finally divided,        entrained particulate);    -   3) that the gas be easily supplied to the system and easily and        rapidly turned on and off;    -   4) that it make a secondary film that will not react with the        primary film at processing or anticipated use temperatures; and    -   5) that it provide an effective barrier to the epitaxial growth        of the tungsten crystallites from one primary film to the next        primary film.

The use of a carbon or boron-containing reactant gases for the additivegas to produce the secondary films satisfies the above requirements.Other compositions for secondary films, such as compositions containingoxygen, are also within the scope of the invention. From an engineeringpoint-of-view, this simple addition of another gas to form the secondaryfilm makes this required gas flow switching to form the very thin layersreasonably simple. It would be much more complicated and createsignificant flow disruption, if the primary film precursor gas wouldalso have to be switched.

In an alternate embodiment, a different kind of precursor can be usedfor the secondary film deposition. For example, silicon carbide might bea useful composition for the secondary film. It is a hard material,making it useful for tool and wear applications for the alloys of theinvention. It has a very acceptable match in coefficient of thermalexpansion with tungsten or molybdenum. It should be effective ininterrupting the epitaxial growth of the primary film grains; and, itwould appear to be possible to inject precursor gases for it without theneed of switching off the flow of the primary film precursor gas. Inaddition, it should be acceptably inert to the primary film metals atboth the processing and contemplated use temperatures.

As noted above, the control of the thickness of the secondary film isimportant. Referring again to the first illustrative embodiment, afterthe selected duration for the flow of the additional gas and theachievement of the desired thickness of the secondary layer, theadditional gas is turned off. This duration of flow for the gas of thesecondary films is shorter than that for the primary films for mostuseful product applications. For very hard coatings or very hardfree-standing bodies, the secondary films may be as thick as, or eventhicker than, the primary films.

There would appear to be a minimum thickness for the secondary film. Anestimate of this thickness can be made, as follows: The atomic diameterof tungsten is about 0.3 nm. A grain boundary is necessary to interruptthe epitaxial growth of the primary film grains. It is taught in thewidely accepted text, Dieter, George E., “Mechanical Metallurgy” fromthe 3^(rd) Ed., McGraw Hill, 1986; that grain boundaries are “a regionof disturbed lattice only a few atomic diameters wide”. This wouldindicate that it would require a secondary film thickness in the orderof 1 nm. to effect the necessary interruption.

Even if thickness of the secondary films is greater than that of theprimary films, the secondary films must be very thin to eliminate thepossibility of large flaws in this material of little of no ductility.This obviously implies that, if the secondary films are to be thickerthan the primary films, the primary films must be exceeding thin.

The ratios of these thicknesses are used to control the desired hardnessof the alloy in addition to the hardness enhancement due to the extremegrain refinement.

After a secondary film is formed, the tungsten halide continues to flow,being decomposed to form another primary (tungsten) layer, after whichthe additional gas is admitted, again, to form an additional secondarylayer and to interrupt growth of the tungsten crystallites again. Theprocess may be repeated forming a multiplicity of very thin films, whichare highly adherent one to the other, until the desired thickness of thecoating or body is attained.

By using this gas-switching technique, microcrystalline material ofgreat constituency and high repeatability is able to be accomplished, asshown in FIGS. 3 and 4. Whichever technique of the present invention isused, the thickness of each of the primary films can be maintainedwithin a 2:1 ratio of any other primary film and the thickness of eachof the secondary films can be maintained within a 2:1 ratio of any othersecondary film. When using the gas-switching technique, the actuationspeed of the switches should be as rapid as possible, as, for example,by the highest-speed-available, solenoid-actuated valves to assure thegrowth of abrupt heterojunctions.

In addition to serving simply to interrupt the growth of thecrystallites of the primary film, the secondary films must be ofsufficient thickness to prevent the continuation of the epitaxial growthof the primary film crystallites in the next primary film, i.e., toprevent the continuation of the epitaxy across the secondary film. Inthe research of Chu and Barnett, referenced above, such interruption isnot effected and a totally different mechanism of ensuring thenanostructure behavior is employed. In the method of the presentinvention, physical vapor deposition could also be effectively used forboth the primary film and the secondary film. The required thickness forprevention of the continuation of epitaxy depends on the difference inthe crystal dimensions and type of the crystallites in the secondaryfilm as compared to those in the primary film.

As noted above, however, whereas the secondary film must be sufficientlythick to interrupt the epitaxy of the primary film crystallites, itcannot be too thick. In one embodiment, the secondary film is thinnerthan the primary film. In an alternate embodiment, the secondary filmmay be thicker than the primary film but less than approximately 400 nm.For the invention to be most effective in creating the unusual physicalproperties in the products made therefrom, the secondary film should notexceed a certain thickness. This thickness is approximately equal to orless than the critical size of defects which would degrade the fracturetoughness of the bodies.

The unique crystal habit has additional advantages:

-   -   a. the alloys are more resistant to recrystallization than        structures of tungsten in which the grain refinement is effected        by mechanical working, and,    -   b. when recrystallization does occur a fine dispersion of        tungsten and tungsten carbide results that maintains much of the        high strength characteristic of the alloy prior to heat        treatment.

It should be clear to one skilled in the art that, in an illustrativeembodiment, the tungsten layer might be formed from a mixture of metalhalides to form a solid solution tungsten alloy suitable for the primarylayer.

Alternatively, a small amount of carbon or boron could be added to theprimary layers; not enough to form a compound like tungsten carbide ortungsten boride, but enough to effect some solid solution strengthening.Such solid solution alloys will tend to improve both the roomtemperature and high temperature strength and hardness of the alloys.

A number of useful alloys of tungsten with other refractory metals areknown. Such alloy additions are not necessary, however. Good propertiescan be attained using pure Group VIB metals, such as tungsten, for theprimary films in the method of the invention. The Group VIB metals arehigh-melting. This makes the alloys useful for applications involvinghigh use temperatures or for tool and wear applications wherein heat isgenerated. It follows, therefore, that, for most useful applications,the material of the secondary film should also be high melting.

Each of the primary and secondary layers is, as noted, a quite thin,microcrystalline layer. Useful dimensions are on the order of 10-1000 nmfor the primary layer, and <1-400 nm for the secondary layer. Forexample, 10 nm for the primary layer can be accomplished by running theprecursor gas for that layer for approximately 0.5 seconds. Thedeposition of a secondary layer having a thickness of less than 1 nmwould require running the additive gas for less than 0.1 seconds. Thereis no intrinsic limitation to the total thickness of the depositedalloy, either for a coating or free-standing body. Overall depositionrates can be as high as 0.1 to 1.0 mm per hour. Massive deposits simplyrequire extended run durations. Acceptable deposition rates for thincoatings can be much lower.

In one embodiment, the invention is primarily directed toward theproduction of thick coatings or free-standing parts. Either of these areparts where external finishing of the deposited material is usuallyindicated. It should be explained that from a deposition point-of-viewthere is really no difference between a massive coating (e.g., one ofabout 0.1 mm to several mm in thickness) and a free-standing part. Theonly difference is that in the case of the massive coating the depositis made to be adherent to the substrate and is left on the substrate;whereas, in the case of the free-standing part it is removed from thesubstrate, either mechanically, by contaminating the substrate surfaceand thereby assuring that it is not adherent, or by liquid dissolutionof an inexpensive substrate/mandrel, or even by reacting the mandrelwith an active gas to completely remove it and form the free-standingpart. Whether it be a massive coating or a free standing part, theproperties of the deposit largely determine the bulk performance of thefinished part.

Even thin coatings of the material of the invention, on the order of 4-8μm, although they do not need the improved mechanical integrity of theimproved material of the invention, will, under certain circumstances,have their performance enhanced by the superior properties of strengthand toughness and hardness.

A method of the invention allows the synthesis of tungsten alloyscovering a wide range of selected compositions from alloys of very nearpure tungsten designed to have very high strength and toughness withmodest hardness, e.g., an HV of 7 GPa, at one extreme (essentiallyproviding superior properties for use in applications where conventionalpure tungsten would be used) to alloys of much greater hardness, e.g.,an HV of 20 GPa or more, having lesser but still relatively highstrength and superior toughness over materials of equivalent hardnessand good wear resistance. The properties are dependent on the size ofthe crystallites (which is essentially determined by the primary filmthickness) and the relative amounts of the primary and secondary layers.

The primary film is sufficiently thin to control its crystallites to asufficiently small size to provide a high yield strength in accordancewith Hall-Petch. On the other hand, the primary film may be sufficientlythick to provide the necessary plasticity in the body.

The secondary film is sufficiently thick to prevent epitaxial growth ofthe crystallites of a preceding primary film across the secondary filminto the following primary film. It is important to note that otherinvestigators have not become aware of the importance of thisrequirement. The secondary film is sufficiently thin to prevent theincorporation of flaws of critical size within the film. Both the sizeof the grains and the ratio of the thicknesses of the primary tosecondary film provides control of the hardness of the alloy.

Regardless of how excellent the bulk properties of the inventionmaterial are, the overall behavior can be further improved by addressingthe surface behavior. Simplistically, it can be stated that to achievethe maximum performance from the invention material, the surface issufficiently free of flaws so that the body does not fail by surfacedefects at a stress level less than it would survive based on its bulkproperties.

There is nothing unique about a discovery that the strength of materialsis improved by eliminating surface flaws. With the grain refinementwhich is to be found in conventional materials, the bulk properties areseldom good enough so that the performance of a body can be limited bythe superficial properties. What has been discovered in this invention,though, is that when a body is made of a multiplicity of very thin filmsby CVD in such a way that epitaxial growth is eliminated, the entirebody is virtually totally free of strength limiting defects. When thisoccurs, the crack initiation stress is raised greatly, but the criticalflaw length is reduced greatly and approaches the size of the very smallcrystallites.

This means that such materials are potentially very strong, but thatthis strength is more effectively realized if the surface finish is verygood. Cahoon's metric of σ_(ys ≈)⅓ HV is observed to be strictlyfollowed for the subject invention at HV of 13 GPa if the surface finishis nearly perfect, i.e., of a roughness of not much greater, if greaterat all, than the grain size of the crystallites in the body (forpractical purposes, the size of the grains in the primary film).Non-trivial plastic deformation for a material of such high hardness isalso observed. It is believed that higher ultimate strengths as well asplastic deformation would also be observed in materials with hardnesseven greater than 13 GPa if the surface defects were to be reduced toclose to or less than the grain size of the primary film, i.e., ˜100 or200 nm. Such good finishes can be achieved by grinding with a veryfine-grit wheel; but they have been more easily achieved, with lessdanger of causing surface flaws, by electro-polishing. Similar finishingmethods like electrochemical grinding or chemical mechanical finishing,should be as effective as electro-polishing.

It should be noted that tungsten was selected for the investigations inconnection with which an illustrative embodiment of the invention wasmade because of its high theoretical properties of strength, stiffness,and strong bonding as compared with other metallic materials.

Originally it was the intent of these investigations to make a hardmetal composition from a mixture of tungsten carbide to provide thehardness and a sufficient amount of pure tungsten metal to providesufficient strength and toughness. To be competitive with commercialcemented tungsten carbide products, the hardness could be selected inthe range from about an HV of 12 GPa to about 22.

Based on the well-accepted Rule of Mixtures and the known hardness of —Wand WC, this would indicate a composition between about 40% WC/60% W andabout 90% WC/10% W. Of course, some hardness enhancement would beexpected due to extreme grain refinement so that the need for such highconcentrations of WC was not really anticipated.

The current invention involves a desirable low fraction of a brittlehard metal compound and a resulting high fraction of nominally pure,softer metal. For even the hardest alloys of the invention the totalvolume of hard tungsten compounds (in the secondary films) does notapproach 100%. At as little as 5% of carbide or boride, it might beexpected, based on the Rule of Mixtures, that these hard species wouldmake the alloy mixture having an HV of about 5 GPa. In fact, when theprimary grains are refined sufficiently alloys of over 20 GPa result. 20GPa is as hard as cemented carbide tool materials. It is apparent tothose of ordinary skill that much hardness enhancement has beenaccomplished from grain refinement in alloys per Hall-Petch. However,the inventors are not aware of its having been accomplished before.Using the method of the invention, very consistent enhancement of bothstrength and hardness is achieved without loss of toughness in thesealloys which have an unusually high concentration of metal (with anunusually small amount of a hard brittle phase).

The composition of the secondary film also has some influence on thehardness of the alloy, particularly in those alloys having a high ratioof secondary film thickness to primary film thickness. The method of theinvention assures each individual primary film can be of the samethickness as any other. Likewise, it allows each individual secondaryfilm the same as any other in the deposit.

In contrast, the method of the '345 patent did not allow for suchnecessary control. In fact, this lack of ability to effect precisecontrol of the layer thicknesses is the probable cause of the poorreproducibility of this former method.

The method of the present invention also allows these primary andsecondary films to be controlled to very small thickness dimensions,even into the nanometer range. Very high strengths, which not beenreported by other nanotechnology investigators, have been achieved.

These increases in hardness may be effected simply by reducing thethickness of the primary films while maintaining the secondary films atthe minimum required for their epitaxial growth interruption function,or by altering both the ratio of primary to secondary and also effectingan increase in the thickness of the secondary film. The utility ofeither of these two methods may be determined experimentally for anygiven product application.

It should be apparent to those skilled in the art that there can be anadvantage in programming a purposeful change in the layer thicknessesthrough the thickness of the deposit. This would allow, for example, onesurface of a free-standing body to be very hard, and wear-resistant, andthe other surface to be very tough and resistant to crack initiation andstructural failure. The method of the invention allows for precisecontrol of primary and secondary film layers of controlled varyingthicknesses, or control of the ratios of primary to secondary filmthicknesses to effect such a purposeful change.

Tungsten, of course, is a metal of great stiffness, about 400,000 GPa.This turns out to be important as far as the strength of the material ofthe invention.

The matter of the importance of the size of internal or surface flaws inthis nanostructural material was discussed above.

It is generally accepted, based on the early work of Griffith, that fora given internal or external flaw (crack) the brittle fracture stressvaries as the square root of the Young's modulus.

That the method of the invention truly minimizes the size of the flawsis very important, but, in addition, tungsten is a preferred commonmaterial upon which the material of the invention can be based.

There is another reason why tungsten is a preferred major species forthe alloy of the invention. As taught by dislocation theory, microcracksnucleated at the head of a dislocation pile-up subjected to a shearstress, can cause brittle fracture if the applied stress can propagatethe microcrack. This failure stress is inversely proportional to thesquare root of the microcrack flaw size and directly proportional to theshear modulus. Therefore a material like tungsten with a higher Young'smodulus (shear modulus is directly proportional to Young's modulus) isto be preferred among the common metals from both a Griffith flawviewpoint as well as from dislocation theory. In addition, the bodycentered cubic crystal structure of tungsten, like iron, demonstratesthe highest Hall-Petch sensitivity; i.e. a reduction in grain sizeresults in a greater strength and hardness increase than with othercrystalline forms. It is believed that other Group VIB transitionmetals, such as molybdenum and chromium, will exhibit similarcharacteristics and are also within the scope of the present invention.

These effects have been verified experimentally. We have clearlyestablished that the flexural strength of the alloy exceeds that ofcommercially available alloys which are in the range of 1800 MPa.Flexural strengths of over 5100 MPa have been achieved. Such strengthshave been attained even for materials with hardness in the useful rangeof some quite hard cemented carbide grades, c.a., HV of >20 GPaSignificantly higher strengths are anticipated to be demonstrated infurther development, achieving a higher percentage of the calculatedtheoretical maximum for tungsten of over 40,000 MPa.

In addition to flexural strengths, there are other properties of theinvention which are important. A comparison of the resilience andtoughness of the alloys of the invention versus other materials and,particularly hard alloys, reveals some important data.

Again, according to Dieter, Ibid.:

-   -   The ability of a material to absorb energy when deformed        elastically and to return it when unloaded is called        “resilience”. This is usually measured by the modulus of        resilience, which is the strain energy per unit volume required        to stress the material from zero stress to the yield stress.

Also from Dieter:

-   -   The toughness of a material is its ability to absorb energy in        the plastic range. . . . Toughness is a commonly used concept        which is difficult to pin down and define. One way of looking at        toughness is to consider that it is the total area under the        stress/strain curve.

It can be seen that for hard, strong materials whose yield strength isvery close to its ultimate strength, like the cemented carbides, thereis very little difference between Modulus of Resilience and Toughnessand values in consistent units would be practically identical. Such isnot the case with more ductile materials like the tool steels. In fact,the following comments on tool steels can be found in the “MetalsHandbook” 1948 Edition, published by the American Society for Metals:

-   -   In tool steels, the concept of toughness is best expressed as        the ability to resist breaking rather than the ability to deform        plastically before breaking, since most tools must be rigid        articles that do not deform in service.

Additionally to the point, some values for the Modulus of Resilience,UR, and Toughness are shown below (determined using flexural strength in3-point bending): Modulus of Resilience Fracture Energy MaterialJoules/m³ Joules/m³ Commercial tungsten  8.2 × 10⁶   14.8 × 10⁶ (asworked, unrecrystallized) Modern, High-Strength, 10.9 × 10⁶ ≈10.9 × 10⁶Micro-grain, Cemented Carbide (˜20 GPa, HV) Tough Alloy of Invention10.6 × 10⁶   32.8 × 10⁶ (7.8 GPa HV) Hard Alloy of Invention 24.2 × 10⁶≈24.2 × 10⁶ (20 GPa HV)

The “tough alloy” is more like a tool steel, having hardness like thatof such steel. It has a high toughness as compared to the most moderncarbides with about the same Modulus of Resilience. The “hard alloy”(about the same hardness as the cemented carbide) has both higherresilience and toughness than the carbide.

It has been determined that the alloys of the invention have anotherdesirable property unrelated to the above. Surprisingly, even thoughrecrystallization is observed after a one hour exposure at 1500° C., theresulting fine dispersion of tungsten and tungsten carbide maintains alarge fraction of the strength of the original alloy. The tough alloycould be used commercially in place of wrought tungsten, tungsten heavyalloys, and CVD based tungsten tubing and other fabricated shapes. Theimproved properties of strength and toughness can result in more robust,higher-performing products.

The high-strength, hard alloys of the present invention could be usedcommercially in place of tool steels, cast hard nickel or cobalt-basedalloys or cemented carbides. They could be used for machining metals andnonmetallics. Furthermore, with the achievable very smooth surfaces,very keen cutting edges can be made. They can also be used for wearparts for machinery such as bearings and seals, and for wear-resistantnozzles such as those used in metal and nonmetallic cutting andfinishing. They could be particularly adaptable to a wide variety ofmining tools such as those used in coal mining, hard rock mining and thedrilling of oil wells. Some of the best applications could be found invalve trim, particularly those involving both corrosion and erosion andin the lining of cylinders and rotary pumps.

To assist in an understanding of the invention, certain examples oftechniques and products are given.

Due to the very small scale of each individual primary and secondaryfilm thickness, direct measurement requires careful polishing andetching of metallurgical cross-sections and the use of very highmagnification (20,000×-100,000×), such as obtainable with a highresolution scanning electron microscope or a transmission electronmicroscope, to directly observe the film thicknesses. To minimize thetime and expense of such measurements relating to this invention, auseful indirect measurement of film thickness was adopted and termed“couplet thickness”. This measurement can be defined as the totalthickness of one primary film and one secondary film. Couplet thicknesswas calculated as follows:d=D*(θ_(p)+θ_(s))/θ_(T)

-   -   where d is the couplet thickness, D the total deposit thickness        (as measured after the run), θ_(p), θ_(s), and θ_(T) are the        duration of a primary film deposit, the duration of a secondary        film deposit, and total run duration, respectively

This parameter tracked the true combined thicknesses of a single primaryand a single secondary film, if, of course, as was the objective of theresearch, all primary films were of the same thickness and all secondaryfilms were of the same thickness. Direct measurement at highmagnification of etched metallographic specimens revealed that thisconsistency was fairly regularly experienced; except at the verybeginning of all runs using a hot wall reactor, where a period of timewas necessary before a stable heat balance among the hot-wall heatsource, the reactant gases, and the work piece was achieved. Inproduction equipment, such an unstable period can be avoided by oneskilled in the art through a simple programming of the furnace energyinput. The issue can be even more easily ameliorated if a reactor usinginternal or induction heating of the workpieces is used. This same kindof programming is one of the ways a purposeful variation of thehardness, strength, and toughness properties can be effected if such isrequired for certain tool and wear applications.

The ratio of the thickness of primary film and that of the secondaryfilm in the couplet was determined in one of two ways. In the first, thedeposition rates for the primary film deposits and the secondary filmdeposits were measured directly from previous calibration runs utilizingthe same gas switching timing for either the metal or metal compoundcycle, which allows the other film thickness to be indirectlycalculated. Additionally the primary film thickness can fairlyaccurately be determined when it is known that the secondary filmthickness is fairly small in comparison. Using these data, it waspossible to dead reckon the metal and metal compound fraction of thecouplet.

The second used the alloy composition data of the metal compoundfraction in the metal as defined by x-ray diffraction. Knowing thefractional relation, it was possible to determine the primary andsecondary film thicknesses of the couplet by incorporating the densityof the component materials. Glow discharge mass spectroscopy data canresult in a more accurate indirect determination of the primary andsecondary film thicknesses if the composition of the metal compound isknown (e.g. WC or W₂C as determined previously by x-ray diffraction).Alternatively the same answers can indirectly be obtained more quicklyand simply by determining the density by Archimedes method.

For all of the following Examples (Except for a different parttemperature in Example 1), the following were held constant. Depositionsubstrates were made of 1 mm diameter molybdenum wire, 16 cm. long. Thewire was racked vertically in a uniform array. Number of samples waseither 4 or 6. Tungsten chloride was made in situ by passing meteredchlorine gas over a bed of tungsten chips, heated to 800° C.Representative atom ratio, Cl/W, of the effluent gas was 4.0. Allpropane flow additions were at a C/W ratio of 2. Ratios for secondaryfilms using other precursors are indicated. All runs included a hydrogenreductant in excess of the stoichiometric ratios to chlorine. Reactordeposition temperature was 900°. These are referred to below as“standard conditions”.

EXAMPLE 1

Samples of tungsten were made by well-known, conventional, CVDprocessing methods. Conditions were as above, except that the parttemperature was held at 800° C. to attempt to prevent undue coarseningof the deposits. Tungsten fabricated at higher temperatures possessedpoorer mechanical properties due to their coarse nature. Resultantspecimens were examined microscopically and the epitaxial growth of thetungsten crystallites shown. The sample demonstrated a strength of 810MPa and a fracture toughness of 1.95×10⁶ Joules/m³. Since there was nomeasurable plastic deformation, the indicated specific resilience wasalso 1.95×10⁶ Joules/m³. An additional sample was electro-polished andtested. This showed a strength of 938 MPa. These were very acceptablestrengths as compared with typical reported data from CVD parts, ortungsten parts made by other means of consolidation. These tests servedas a base line against which to compare the materials of the invention.

EXAMPLE 2

Two tests were run to determine the minimum thickness of the secondaryfilms which would effect interruption of the epitaxial growth of thetungsten metal grains in the primary films.

The test using a primary film flow duration of 3 seconds and a secondaryfilm flow duration of 0.1 seconds indicated by visual examination of thetopography and fracture surfaces of the sample that interruption hadoccurred. The sample, after electropolishing, had a measurable strengthof 2238 MPa, confirming this indication. Estimating from depositionrates experienced in other tests, the secondary film thickness wascalculated to be approximately 2 nm.

The test was repeated using a secondary film flow duration of 0.05seconds. Similar visual examination of the sample topography andfracture surfaces suggested that the epitaxial growth had not beeninterrupted on most of the sample. The upstream 10% of length showed asmooth enough surface to suggest some interruption on this portion ofthe sample. In addition, microscopic examination of a section of a lowerportion of the sample showed clearly the presence of columnar grainsresembling ordinary CVD tungsten for the outermost ⅔ of the deposit, buta featureless deposit for the innermost ⅓.

Strength was measured at 379 MPa; actually lower than the base linematerial.

These tests indicated that the duration of 0.05 seconds secondary filmprecursor flow was not quite long enough and was marginally ineffective,in all probability because it produced too thin a secondary film. Thislatter secondary film thickness would have been estimated to be about1.0 nm based on extrapolation of deposition rate data from earlier runs.This is in substantial agreement with the hypothesis offered above,having to do with the minimum thickness of secondary film required.

EXAMPLE 3

A test was conducted using the standard conditions described above.Tungsten chloride flowed continuously. Injections of propane gas toproduce tungsten carbide were made, periodically. The primary filmdeposition flow duration was 3.0 seconds. The secondary film flowduration was 0.2 seconds, which is close to the minimum secondary filmduration of deposition which had proved effective in interrupting theepitaxial growth of the primary film tungsten grains in the earliertests, as shown in EXAMPLE 2. Using the dead-reckoning method, thethickness of the primary film was calculated to be 240 nm and that ofthe secondary film, 4 nm.

The ultimate strength was measured at 3503 MPa and the yield strength at2566 MPa after electropolishing. The hardness was 8.0 GPa.

EXAMPLE 4

A run similar to that of EXAMPLE 3 was made. The same standardconditions were used. In this case, however, the primary film depositionflow, without propane injection, was shorter, for 1 second duration. Thepropane gas injection was for the same 0.2 seconds. The cycle wasrepeated for 5 hours.

Again using dead-reckoning, the thicknesses of the primary and secondaryfilms were calculated to be 61 nm and 4 nm, respectively. Specimens weretested in three point bending after electropolishing. The ultimatestrength was measured at 4440 MPa. The yield strength was 4172 MPa. Thehardness (HV₁₀₀₀) was 12.0 GPa. The results were as expected. Theextreme grain refinement produced the high strength and the higherhardness than that shown in the run of EXAMPLE 3.

From both of these runs the difference provided by grain refinementproduced by the method of the invention can be seen. These samples arebetween 4 and 5 times the strength of the conventional CVD samples ofEXAMPLE 1. The specific modulus of resilience and the fracture energyare equivalently higher.

EXAMPLE 5

A run similar to that of EXAMPLE 3 was used to demonstrate the utilityof a secondary film of different composition. Boron trichloride flow ata B/W atom ratio of 2.5 was used as the secondary film depositprecursor. Primary film flow duration was 10 seconds; secondary filmflow duration was 0.5 seconds. As was expected, the grains in the samplerefinement were not extraordinarily fine, but still quite small at400-500 nm. The resultant sample was measured at an ultimate strength of3324 MPa after electropolishing. This Example demonstrates theequivalence of tungsten boride with tungsten carbide in interrupting theepitaxial growth of the tungsten grains. Note again the high strengthrelative to the baseline material.

EXAMPLE 6

A test was conducted using the standard conditions described above. Thistest was similar to the tests of EXAMPLE 3. Tungsten chloride flowedcontinuously. Injections of propane gas to produce tungsten carbide weremade, periodically. The primary film deposition flow, without propaneinjection, was for 2.3 seconds duration. The secondary film depositionwas for a considerable longer time than in EXAMPLE 3, 6.0 seconds. Theresultant material was, as expected, much harder than the samples ofEXAMPLE 3. The hardness, measured with a 1000 g. load, averaged 21.6GPa. Specimens were tested in three point bending afterelectropolishing.

The ultimate strength was measured at 3434 MPa. No yielding wasobserved. Metallographic examination of an etched section of the sampleat 20,000× in a scanning electron microscope showed the averagethickness of the primary layers to be 232 nm and that of the secondarylayers to be 372 nm. This example demonstrates that the hardness will beincreased by increasing the secondary film thicknesses without adisqualifying reduction in strength and toughness.

EXAMPLE 7

A test was conducted using the standard conditions. This test was alsosimilar to the tests of EXAMPLES 3 and 4 except that a much longerprimary film flow duration was used, 12 seconds. A short time, 0.2seconds, was used for the secondary film flow duration of propane gas toproduce tungsten carbide. These conditions produced the expectedmaterial which was somewhat softer than the samples of EXAMPLES 3 and 4because they had a smaller proportion of carbide, and because they werenot as fine-grained as the either the EXAMPLE 3 or 4 material. Thehardness was 5.8 GPa. The grain size was dead reckoned to beapproximately 1350 nm (1.35 μm). The strength, 3090 MPa, was still muchhigher than the base-line material, however.

EXAMPLE 8

A test was run with the objective of producing fine primary filmthicknesses with a higher ratio of carbide to metal; to demonstrate thathigh hardness and high strength could be provided in the same sample.The same standard conditions were used. Primary film deposit durationwas 0.8 seconds. Secondary film deposit duration was 0.5 seconds. Theaverage couplet thickness was dead-reckoned at 34 nm. Using W and W₂Cdeposition rates from earlier runs as a basis for calculation, filmthicknesses were estimated as 19 nm for the primary and 15 nm for thesecondary. The resulting samples had an average hardness of 20.2 GPa andan average strength of 4455 MPa. One sample showed a strength of 5172MPa.

EXAMPLE 9

A run was made using the standard conditions, above. The timing was 3.0seconds for the primary gas flow and 0.2 seconds for the secondary gasflow. The resultant specific gravity for the deposited material was19.14 g/cc and the hardness 8.8 GPa. The as-deposited ultimate flexuralstrength was 1959 MPa. Note that this strength is more than twice thatof the baseline material. A companion sample, ground with a 120 gritdiamond wheel (116 μm), was tested at a strength of 1910 MPa. Anidentical sample from the same run was ground and electropolished to aspecular finish (measured at an R_(a) of 200 nm by atomic forcemicroscopy). It tested at a much improved, flexural strength of 3090MPa.

Another run was made using the same conditions except that the timingwas 2.3 seconds for the primary film and 1.1 seconds for the secondaryfilm. For this sample the specific gravity was 18.85 g/cc and theaverage hardness 14.3 GPa. The as-deposited ultimate flexural strengthwas 1635 MPa. A companion, and apparently identical, sample from thesame run, as-ground with a coarse (120 grit) diamond wheel, (116 μm),demonstrated a strength of 1876 MPa. The strength of another companionsample, around with a much finer, 1000 grit (9.2 μm), wheel having amore compliant binder, again producing a specular finish (measured at anR_(a) of 200 nm.), was tested to a much higher strength of 3641 MPa.

This example demonstrated that the bulk strength of these fine-grainedmaterials is best utilized when an excellent surface finish, to avoidsurface defects, is provided.

EXAMPLE 10

A run was made which was similar to Runs 3 and 4 except that instead ofpropane (C₃H₈) gas for a secondary film deposition precursor, propylene(C₃H₆) was used. This precursor was clearly effective, resulting in asample with a strength of 2414 MPa. It was interesting to note that thecomposition of the secondary film was different from those with thepropane precursor, being WC instead of W₂C. Another run was made usingmethane gas, CH₄, for the secondary film precursor. It was alsosuccessful, producing a sample with a strength of 3319 MPa.

EXAMPLE 11

Two samples from the runs described in EXAMPLE 8 were subjected to avacuum heat treat at 1500° C. for one hour. The first sample which had ahardness of 8.1 GPa, as-deposited, showed a hardness of 4.6 GPa afterthe heat treat. The strength of these samples which had been ground witha 400 grit (37 micron) wheel, and was, therefore, not as strong aselectropolished samples, was reduced by the heat-treat from anas-finished value of 2317 MPa to 1697 MPa, still substantially betterthan the base-line material.

The second sample, which had a hardness of 14.0 GPa, as-deposited,demonstrated a softening to 5.0 GPa after the 1500° C. heat-treat. Thissample, which had been ground with a standard 1000 grit (9.2 micron)wheel and electropolished, had an as-finished strength of 2690 MPa and astrength after heat-treat of 2200 MPa. Metallographic examination ofetched cross-sections revealed that the layered structure ceases toexist, but that the material is composed of a fine grained mixture oftungsten and tungsten carbide. These observations indicate that thematerial is more resistant to recrystallization and grain coarseningthan conventional tungsten materials, and thus maintains a higherstrength after such elevated temperature exposures.

From the foregoing detailed description, it will be evident that thereare a number of changes, adaptations, and modifications of the presentinvention which come within the province of those skilled in the art.The scope of the invention includes any combination of the elements fromthe different species or embodiments disclosed herein. However, it isintended that all such variations not departing from the spirit of theinvention be considered as within the scope thereof.

1. A high strength alloy comprising: a plurality of depositedmicrocrystalline films wherein said microcrystalline films alternatebetween a primary microcrystalline film comprised of a Group VIBtransition metal and a secondary microcrystalline film comprised of ametal, a metal compound, a semi-metal, or a semi-metal compound; whereinsaid secondary microcrystalline film has a crystal habit different fromthe body-centered-cubic habit of the metal of the primarymicrocrystalline film and has limited solubility or reactivity withrespect to said metal at the consolidation and use temperatures of thealloy, and said secondary microcrystalline films are disposed betweensaid primary microcrystalline films; and wherein each of said primarymicrocrystalline films has a thickness of up to 1350 nanometers andwherein each of said secondary microcrystalline films has a thicknesssufficient to arrest the growth of the crystallites of the primary filmsand prevent epitaxial growth between adjoining primary microcrystallinefilms.
 2. The alloy of claim 1 wherein said secondary microcrystallinefilms each have a thickness that is less than the thickness of theprimary microcrystalline films.
 3. The alloy of claim 1 wherein saidsecondary microcrystalline films each have a thickness that is greaterthan the thickness of the primary microcrystalline films but less than400 nm.
 4. The alloy of claim 1 wherein said secondary microcrystallinefilm comprises a compound of tungsten.
 5. The alloy of claim 4 whereinsaid compound of tungsten comprises tungsten carbide.
 6. The alloy ofclaim 4 wherein said compound of tungsten comprises tungsten boride 7.The alloy of claim 1 wherein said secondary microcrystalline filmcomprises silicon carbide.
 8. The alloy of claim 1 wherein said alloyfurther comprises a coating or a body having a surface finish (Ra)better than 200 nm.
 9. The alloy of claim 8 wherein said coating or bodyfurther comprises an exterior surface finished by electro-polishing,electro-chemical grinding, or chemical-mechanical finishing.
 10. A highstrength tungsten alloy comprising: a plurality of adherentmicrocrystalline, chemically-vapor-deposited, films wherein saidmicrocrystalline films alternate between a plurality of uniformlydisposed tungsten films and a plurality of uniformly disposed secondaryfilms comprised of a hard metal compound of tungsten; and wherein thethickness of each of the tungsten films does not exceed 1350 nanometers;and wherein said secondary films have a thickness sufficient to arrestthe growth of the primary film crystallites and to prevent epitaxialgrowth between said tungsten films.
 11. The alloy of claim 10 whereinsaid secondary films each have a thickness that is less than thethickness of the tungsten films.
 12. The alloy of claim 10 wherein saidsecondary films each have a thickness that is greater than the thicknessof the tungsten films but less than 400 nm.
 13. The alloy of claim 10wherein the thickness of any tungsten film does not differ in thicknessfrom any of the other tungsten films in the adherent array by more than2:1 and wherein the thickness of any secondary film does not differ inthickness from any of the other secondary films in the adherent array bymore than 2:1.
 14. The alloy of claim 10 wherein said tungsten filmseach have a thickness between 10 and 1000 nanometers.
 15. The alloy ofclaim 10 wherein the thickness of the primary films do not exceed 100nanometers each.
 16. The alloy of claim 10 wherein the thickness of thesecondary films is not greater than 20 nm.
 17. The alloy of claim 10wherein the thickness of the secondary films is not less than 1 nm. 18.The alloy of claim 10 wherein the flexural strength of the alloy exceeds1800 MPa.
 19. The alloy of claim 10 wherein the flexural strength of thealloy exceeds 5100 MPa.
 20. The alloy of claim 10 wherein the alloy hasa hardness (HV) of about 7.8 GPa and a Modulus of Resilience, determinedusing flexural strength in 3-point bending, of over 10×10⁶ Joules/m³.21. The alloy of claim 10 wherein the alloy has a hardness (HV) of about7.8 GPa and a Fracture Energy, determined using flexural strength in3-point bending, over 20×10⁶ Joules/m³.
 22. The alloy of claim 10wherein the alloy has a hardness (HV) of over 15 GPa and a Modulus ofResilience, determined using flexural strength in 3-point bending, over15×10⁶ Joules/m³.
 23. The alloy of claim 10 wherein the alloy has ahardness (HV) of over 15 GPa and a Fracture Energy, determined usingflexural strength in 3-point bending, over 15×10⁶ Joules/m³.
 24. Thealloy of claim 10 wherein the alloy has a hardness (HV) of over 10 GPaand displays nonlinear stress-strain behavior, and such nonlinearbehavior increases at elevated temperatures.
 25. The alloy of claim 10wherein said alloy has been heat annealed at a temperature of 1500° C.for one hour in an inert environment causing a reduction of hardness toless than 6 GPa but a maintenance of strength of at least 1350 MPa. 26.The alloy of claim 10 wherein said alloy further comprises a coating ora body having a surface finish (Ra) better than 250 nm.
 27. The alloy ofclaim 10 wherein said coating or body further comprises an exteriorsurface finished by electro-polishing, electrochemical grinding, orchemical-mechanical finishing.
 28. A method of making a high strengthalloy, the method comprising the steps of: depositing a Group VIBtransition metal in a primary film by physical vapor deposition orchemical vapor deposition on a substrate; using the same depositionmethod to deposit an adherent film of a metal or a semi-metal compoundhaving a crystal habit different from the body-centered-cubic habit ofthe metal of the primary film and having limited solubility orreactivity with respect to said body-centered-cubic metal at thedeposition and use temperatures of the alloy; wherein the deposited filmof said metal or semi-metal compound is of a thickness sufficient toarrest the growth of the crystallites of the primary film and preventepitaxial growth between adjoining primary microcrystalline films; andrepeating the alternate deposition of the adherent Group VIB transitionmetal films and the adherent metal or semi-metal compound films until aplurality of such alternate films is made to the required thickness ofthe coating or body.
 29. The method of claim 28 wherein the primary filmis deposited so that it has a thickness not greater than 1350 nm inthickness.
 30. The method of claim 28 wherein the deposition methodcomprises chemical vapor deposition.
 31. The method of claim 28 whereinthe Group VIB transition metal comprises tungsten.
 32. The method ofclaim 28 wherein the metal or semi-metal compound comprises tungstencarbide.
 33. The method of claim 28 wherein the metal or semi-metalcompound comprises tungsten boride. 34-48. (canceled)